Nano-structured materials have emerged as very important engineering materials in recent years. They are made of crystals with sizes below 100 nm. In these materials 50% of the actual volume consists of grain boundaries. Because of these large numbers of crystalline interfaces an important fraction of the materials have a disordered microstructure with no short-range order. As a result, the nano structured materials exhibit physical and chemical properties different from those usually found in coarse grain crystalline materials. While a significant number of nano-structured materials have been developed in recent years, the application of nano technology in bulk structural materials like steel and cast iron has been rather limited.
The term austempered ductile iron (ADI) describes a family of materials whose properties can be varied over a wide range by the correct choice of heat treatment variables and chemical composition. ADI is an alloyed and heat-treated ductile (or nodular) cast iron. ADI has become a major engineering material due to its excellent properties; these include high strength with good ductility, high wear resistance, good fatigue strength, and fracture toughness. These properties are a result of the development of a unique acicular matrix structure that consists of high carbon austenite (γHC) and ferrite (α) with graphite nodules dispersed in it. Compared to conventional ductile iron, ADI has nickel, copper and molybdenum added to increase its heat treatability; i.e. to delay the austenite decomposition to pearlite and ferrite upon cooling. Proper austempering heat treatment avoids the formation of unwanted microstructural constituents (such as martensite, carbide and pearlite). ADI has low production costs due to its good castability, excellent machinability and shorter heat treatment processing cycles. Because of these properties, it has been used in a wide variety of applications, including gears, crankshafts, locomotive wheels, connecting rods, and brake shoes etc.
The development of ADI involves two major processing steps. The first step is the melting and casting of ductile cast iron that has been specifically alloyed with elements such as Ni, Cu, and Mo. The second processing step is the heat treatment. The casting is heated to, and held at, temperatures ranging between 815-927•° C. (1500-1700•° F.) for one to two hours. This allows the microstructure to become fully austenitic (γ). After austenitizing, the alloy is quenched in a molten salt bath to an austempering temperature ranging between 260-400•° C. (500-750•° F.). The casting is kept at temperature for two to four hours; following this, it is air cooled to room temperature. FIG. 1 provides a schematic of a prior art process for forming ADI. As indicated by the line from A to B (A-B), an ductile cast iron is heated to a temperature at which conversion to austenite occurs. The ductile cast iron is held at this temperature for several hours as indicated by B-C. The ductile cast iron is then quenched to an austempering temperature as indicated by C-D and held at this temperature for two hours (D-E). The alloy is then cooled to room temperature as indicated by D-E.
During austempering, ADI goes through a two-stage phase transformation process. In the first stage, the austenite (γ) decomposes into ferrite (α) and high carbon austenite (γHC):γ•→α+γHC  (Eq. 1).
If the casting is held at the austempering temperature for too long, a second (and undesirable) reaction occurs. In this reaction, the high carbon austenite can further decompose into ferrite and carbide:γHC•→α+ε  (Eq. 2).
In this case, the ε carbide will make the material brittle; therefore, this reaction must be avoided. In general, the optimum combination of tensile strength and ductility is obtained in ADI after the completion of the first reaction but before the onset of the second reaction. The time period between the completion of the first reaction and the onset of the second reaction is called the “process window”. The process window can be enlarged by addition of alloying elements such as Ni, Mo, and Cu.
Proper austempering produces a unique microstructure that consists of high carbon (or transformed) austenite (γHC) and acicular ferrite with graphite nodules dispersed in it. The γHC is present in the form of small “slivers” located between the ferrite needles. The exact morphology of the ferrite phase and the relative amounts of ferrite and γHC can be controlled by austempering temperature and time.
During the austempering of ductile cast iron, acicular ferrite grows from austenite by the nucleation and growth process. As the ferrite grows, the remaining austenite becomes enriched with carbon. The form that this ferrite takes is dependent upon the austempering temperature. When the austempering temperature is in the lower bainitic temperature range i.e, between 232-316•° C. (450-600•° F.), a microstructure consisting of bainitic ferrite (αB), austenite (γHC), and graphite nodules is developed; the bainitic ferrite in this case consists of needle-shaped particles of aggregated ferrite and precipitation of carbide within.
When the ADI is austempered in the upper bainitic temperature range (above 316•° C. (600•° F.)), an acicular structure of carbide-free ferrite (αCF) with a considerable amount of stabilized austenite (γHC) develops. This microstructure is referred to as an Ausferritic (αCF+γHC) microstructure. It contains an interlocking aggregate of fine, randomly oriented, intergranular laths of ferrite with an aspect ratio of 4:1. Because of the high silicon content in ADI, the formation of cementite phase, normally associated with the bainitic reaction in steel, is suppressed in this case. Consequently, the remaining austenite continues to be enriched with carbon as the reaction proceeds. As the austenite becomes enriched with carbon, growth of bainitic ferrite platelets is inhibited, and the reaction is arrested.
Since, the transformation that produces acicular ferrite and γHC is a nucleation and growth process, and the nucleation depends on supercooling, as the austempering temperature decreases, the degree of supercooling increases; therefore, more ferrite is nucleated and the ferrite, as a consequence, becomes finer in nature. Additionally, in the presence of silicon, carbon rejected from the growing ferrite phase during transformation does not form carbides. Instead, the carbon enters into solid solution in the remaining austenite, enriching the carbon content of this austenite. After a certain austenitizing time, the carbon content of the remaining austenite is sufficiently enriched so that its Ms (martensite start) temperature is depressed below room temperature. This results in formation of stable, high carbon austenite (γHC). However, as the austempering temperature decreases, the growth rate of ferrite needles decreases as well. This causes the ferrite and γHC in the matrix to become finer in scale, with a resulting increase in the volume fraction of ferrite.
In contrast, as the austempering temperature increases, the degree of supercooling decreases while the growth rate of ferrite increases. Consequently, the volume fraction of ferrite content decreases, the volume fraction of γHC increases, and both the ferrite and austenite becomes coarser in nature.
Altering the microstructure will alter the resulting mechanical properties in ADI. For example, when austempering is performed at temperatures near 260•° C. (500•° F.), the resulting ADI has a large amount of fine ferrite and γHC in the matrix. Thus, tensile strengths up to 260 Ksi (1600 MPa) with 1 percent elongation and hardness values in excess of 60 Rc are obtainable. Conversely, when austempering is performed at temperatures near 385•° C. (725•° F.), the ferrite and γHC become more coarse and “feathery”; this results in tensile strengths of 120-170 Ksi (800 to 1200 MPa)) and elongations up to 14%.
In most conventional materials, the high-cycle fatigue strength increases as the monotonic yield strength or ultimate tensile strength increases. However, a number of researchers have reported that ADI shows the opposite behavior. These studies found that the fatigue strength of ADI is higher when its yield strength is lower. Thus, the high-cycle fatigue strength of ADI increases with increasing austempering temperature. The higher fatigue strength at higher austempering temperatures (with consequently lower yield strength) is due to the presence of a greater volume fraction of γHC in the matrix. Austenite is a face-centered cubic phase; it has a higher toughness and work hardening rate compared to the body-centered cubic ferrite. Thus, as the amount of austenite in the matrix increases, a higher work hardening rate is present; hence, this leads to high fatigue strength in ADI.
As detailed previously, as the austempering temperature increases with both ferrite and austenite becomes coarser. This coarser γHC is thought to affect the mechanical properties of ADI through increased plasticity in the matrix. One study on transformations in ADI reported that the austenite was found to transform at the crack tips to offset necking instability; this was similar to the results found for lowcarbon transformation induced plasticity (TRIP) and medium carbon forging steels. However, the study observed that this behavior occurred only when the carbon content of the austenite was relatively low.
Previous investigations on ADI with an Ausferritic microstructure have shown that they possess improved fracture toughness. When the fracture toughness is plotted against the austempering temperature, it is found that the fracture toughness initially increases with increasing temperature, reaches a maximum at an intermediate temperature and decreases with further increase in temperature. This is believed to be the result of competitive interplay between the effect of ferrite grain size and the effect of γHC volume fraction. Ferrite has the maximum fracture toughness at the lowest austempering temperature; this is due to the fineness of the grain size developed at the low austempering temperatures. The fracture toughness of the γHC is maximized at the higher austempering temperatures; this is due to the increased volume fraction of γHC (relative to the ferrite). Thus the actual fracture toughness of the ADI is controlled by the “weakest link”; this is the γHC created at low austempering temperatures and the ferrite created at high austempering temperatures.
The relationship between the volume fraction of γHC and the carbon content of the γHC is key to understanding the fracture toughness. This PI has developed an analytical model that is valid for in austempered ductile irons:KIC2=σy(XγCγ)1/2  (Eq. 3)where KIC is the fracture toughness, σy is its yield strength, Xγ is the volume fraction of γHC, and Cγ is the carbon content of the γHC. Other researchers have confirmed the validity of this model. The relationship shown in Equation 3 shows that the fracture toughness of ADI can be maximized by: (a) Increasing its yield strength (σy); and/or (b) Increasing the austenitic carbon content (XγCγ). The yield strength (σy) of ADI depends on the ferritic cell size and volume fraction of austenite. Researchers have shown that the σy depends on width of the ferrite, L, and varies as L−1/2. It has also observed a similar relation between the yield strength of ADI and ferritic cell size. Thus, by producing very fine-scale ferrite and austenite in the matrix, the yield strength of ADI can be optimized. Fine scale ferrite and austenite will also increase the impact strength of ADI.
Increasing the carbon content of austenite will increase the toughness of ADI, as it will result in greater interactions between dislocations and carbon atoms. The carbon content of the transformed austenite (γHC) depends on the carbon content of the initial austenite (γ) as well as austempering time and temperature. During the austempering process, as the ferrite needles grow, the austenite becomes enriched with carbon; this enrichment in carbon content will depend on the austempering time as well as temperature. Thus, if a carbon partitioning mechanism can be developed so that carbon content of austenite will be increased rapidly, then this mechanism will help in reducing the austempering processing time and at the same time will increase the fracture toughness, fatigue strength and yield strength of ADI.
In recent years, significant research has been conducted on the processing of nano-structured materials. Numerous approaches have been investigated, such as alloying, controlled rolling combined with accelerated cooling, plastic deformation and recrystallization (PDR), and repetitive corrugation and straightening (RCS). These techniques have been used to reduce the grain size down to the nanometer scale. However, all these methods have severe problems; many of them produce microporosity and contamination, which results in extremely brittle materials. The published literature indicates nearly all nano-crystal metals have tensile elongation-to-failure values much lower than their conventional counterparts; this is true even for those FCC materials that are very ductile in coarse-grained form.
Further, the literature details this nanostructure material development has focused on steel and nonferrous alloys. Virtually no investigations have been conducted to produce ADI with a nano-scale microstructure.
Accordingly, there is a need for improved methods of making nanostructured ADI having better properties.